1. Field of the Invention
The present invention relates generally to ceramic compositions and, more particularly, to reinforced ceramic compositions, and still more particularly to a novel, erosion-resistant, amorphous or nanocrystalline, ceramic-matrix composition, primarily for fabrication of continuous-fiber-reinforced ceramic composition, distinguishable from sintered powder compositions.
2. Description of the Related Art Including Information Disclosed Under 37 CFR 1.97 and 37 CFR 1.98
The following description of the art related to the present invention refers to a number of publications and references. Discussion of such publications herein is given to provide r more complete background of the scientific principles related to the present invention and is not to be construed as an admission that such publications are necessarily prior art for patentability determination purposes.
The efficient production of continuous-fiber-reinforced ceramics (CFRCs) is desirable for high-temperature aerospace and gas turbine applications. Those materials present an optimal combination of low density (˜30-50% of a metal density) and the strength and toughness imparted by incorporation and embedding of long fibers of carbon or silicon carbide into a ceramic matrix. The ceramic matrices may comprise, for example, without limitation, SiC, SiNC, Si3N4, SiAlON, or the borides, nitrides, or carbides of refractory metals (e.g., Zr, Hf, Ta, W, etc). The CFRCs are a subset of a group of materials known as ceramic matrix composites (CMCs), which comprise a well inter-dispersed, intimate mixture of ceramic and reinforcement filler phases. The filler can be particles, spheroids, whiskers, needles, chards, turnings, filaments, fibers, chopped fibers, fiber thread, fiber cloth, or fabric of metals, polymers, glasses, or ceramics.
Ceramic phases are very strong, monolithic ceramic bodies, which may be single or multiphase compositions. They are, however, prone to catastrophic failure by cracking due to their brittle nature. Long continuous fibers are preferred as the reinforcement vehicle because they provide high ultimate tensile strength as shown by J. A. DiCarlo, H. M. Yun, G. N. Morscher, R. T. Bhalt, in SiC/SiC Composites for 1200° C. and Above; Handbook of Ceramic Composites, N. P. Bansal, ed. Springer Publishing, 2005, Chapter 4. Including long continuous fibers also more effectively mitigates the brittle fracturing of ceramics, a dominant mechanism of the failure of monolithic ceramics, where mechanical failure is governed by the growth of cracks. On a simplistic level, long fibers embedded in ceramics mitigate brittleness in somewhat the same way embedded rebar mitigates the cracking and disintegration of reinforced concrete.
Regardless of the type of reinforcement, all work best when they are fine in size or diameter, closely spaced (i.e., a high numerical density), and preferably geometrically ordered in the body being reinforced. In addition to strength and toughening, long fibers have the advantage that they may be assembled, such as by weaving, braiding, or the lay-up of fabric, into a preform of nearly any shape prior to formation or addition of the ceramic matrix. That process is known as densification or processing of the preform.
Ceramic matrix composites containing 2-D weave fiber reinforcement are known to undergo damage under high tensile load. Such damage can occur first by formation of cracks (transverse to loading direction) in the intertow matrix, between the plies or at yarn intersections within plies. Damage has also been detected in transverse tows (Type 1 and II cracking) typically for strain up to 0.2%, which results in a major decrease in elastic modulus (70%) (J. Lamon “Chemical Vapor Infiltrated SiC/SiC Composite” Handbook of Ceramic Composites, N. P. Bansal, ed. Springer Publishing, 2005, Chapter 3). These cracks initiate at macropores located between plies or next to tow intersections. The big loss of modulus means the load is now carried essentially by the matrix infiltrated longitudinal tows. For deformations >0.2%, transverse microcracks occur within longitudinal tows (stage III), which are arrested by fibers coated with a thin interfacial layer (e.g., ≦1 μm of CVI PyC, BN, or in situ BN+CVI BN) that provides debonding of matrix from fibers and deflection of the cracks. During this stage, the modulus drops by only another 10% as strain approaches 0.6% and matrix damage and debonding from fibers becomes complete (i.e., saturation). The model of strain-induced damage clearly suggests that the elimination of macropores and voids would reduce localization of stresses and initiation of cracking, thereby leading to better toughness.
Several attempts have been made to produce CFRCs with a matrix comprised of a refractory metal ceramic (RMC)—i.e., borides, nitrides, or carbides of a refractory metal—or a mixed phase matrix composed of RMC and SiC. Those attempts include: (1) filament winding and deposition of a slurry containing RMC and SiC powders followed by hot pressing; (2) mixing RMC powder into a polymeric precursor to SiC, which forms SiC during pyrolysis, prior to impregnating fiber or fiber fabric with the powder-precursor mixture during lay-up into a preform followed by pyrolysis and polymer infiltration-pyrolysis cycles; and (3) mixing fiber with a slurry of RMC powder prior to or during fiber lay-up and molding into a preform followed by densification of the preform with SiC. For example, Tang et al. attempted fabrication of a carbon-fiber reinforced ZrB2—SiC matrix composite by chemical-vapor-infiltration (CVI) deposition of the SiC matrix into a carbon fiber-ZrB2 powder preform (see Tang, S., et al, J. Am. Cer. Soc., May 7, 2007). The Tang et al. composite was reported to have a flexural strength of 148 Mpa and a fracture toughness of 5.6 Mpa-m1/2, which compares poorly to the flexural strength of monolithic SiC (205-381 MPa) prepared by chemical vapor deposition. Tang et al. mixed the ZrB2 powder with carbon fibers using a fiber-powder molding technique and then the resulting preform was densified by CVI of SiC matrix. Tang et al.'s materials should not be considered high performance CFRCs, because the ZrB2 phase, which was derived from preformed crystalline powder (average particle size 1.5 μm), contained large grains and was not distributed evenly throughout the fiber perform. Further, the ZrB2 phase was probably not present to any significant amount inside the fiber yarn, (i.e., fiber tow), due to the difficulty of pushing powder into the small (i.e., few micrometers and submicrometer) spaces between the fibrils or filaments of the yarn. Incorporated powder evidently agglomerated and resulted in large (≧10 μm) domains of the ZrB2 phase, and the material had quite low ZrB2 content and low fiber content: ZrB2 content was only 3.9% by weight and fiber volume fraction was only 24.9%. Fiber volume fraction near 40% is desired for high-strength.
A low quality CFRC comprised of SiC fiber reinforced ZrB2 plus 20 vol % SiC has been prepared by a filament winding and slurry deposition technique followed by hot pressing. In that instance, the panel density was much lower than the theoretical density for 35 vol % fiber loading (3.47 vs. 4.60 g/cm3). Fiber volume within the plies was 30%, but thick matrix layers between the plies reduced the overall fiber volume fraction to 20-25% and significant amounts of porosity were seen in the matrix. Based on the apparent density and a rule-of-mixtures calculation, the matrix was about 70% dense, i.e., 30% pores; see S. R. Levine et al. (J. European Ceram. Soc., 22, 2757-2767 (2002)).
Thorough mixing of small RMC and SiC domains (i.e., grains) on a microscopic level (the finer the better) is typically required to obtain optimum mechanical, thermo-structural, and thermo-chemical properties of the composite. R. A. Cutler, “Engineering Properties of Borides;” pp. 787-803 in Ceramics and Glasses, Engineered Materials Handbook, Vol. 4. Edited by S. J. Schneider Jr., ASM International, Materials Park, Ohio, 1991. P. T. B. Shaffer, “Engineering Properties of Carbides;” ibid. pp. 804-11. For example, as reported by A. L. Chamberlain et al. (J. Am. Ceram. Soc., 87(6), 1170-1172 (2004)), the increases in strength and toughness of hot-pressed, monolithic ZrB2 ceramic upon addition of 20 or 30 vol % SiC (from 565 to >1000 MPa strength and 3.5 to 5.3 Mpa-m1/2 fracture toughness) were attributed to a decrease in average grain size. Analysis showed that the grain size decreased from ˜6 μm for ZrB2 alone to ˜3 μm for ZrB2 containing 20 or 30 vol % SiC. The billets contained about 2 vol % WC contaminant for all of the above compositions, as a result of attrition milling the commercial ZrB2 and SiC powders (2 and 0.7 μm particle size, respectively) using cobalt-bonded WC media and a cobalt-bonded WC spindle, so the measured strength for ZrB2 of 565 MPa was significantly higher than strengths reported for phase-pure ZrB2. S. R. Levine et al. (J. European Ceram. Soc., 22, 2757-2767 (2002)) illustrates another example of the importance of thorough intermixing of RMC and SiC domains. The work of the Levine group used fractographic analysis of hot-press sintered ZrB2 containing 20 vol % SiC by, which was composed of equiaxed grains of ZrB2 (6 to 12 μm wide) and elongated grains of SiC (1.5-3 μm thick by 3-11 μm long) to show that fracture during flexural strength testing was often attributed to a number of features. Those features included “large clusters of relatively coarse ZrB2 grains surrounding groups or clusters of large SiC grains.” The Levine, et. al material was virtually pore-free and individual large SiC or ZrB2 particles of width ≧20 μm were not observed. The features often took on a more or less spherical shape indicating the possibility of spherical agglomerates that had not separated from the surrounding matrix during sintering.
Other investigations have reported on the action of SiC as a grain-growth inhibitor in the sintering of TiB2 and ZrB2. See, e.g., R. Telle, L. S. Sigl, and K. Takagi, “Boride-Based Hard Materials;” pp. 803-945 in Handbook of Ceramic Hard Materials. Edited by R. Riedel, wiley-VCH, Weinheim, Germany, 2000. Furthermore, F. Monteverde has reported that addition of 10 vol % ultra-fine SiC particulate was the key factor that enabled both control of the diboride grain growth (average size=3 μm) and the achievement of full density, as well as the enhancement of the strength and oxidation resistance of ZrB2; see, Appl. Phys. A 82, 329-337 (2006), Beneficial effects of an ultra-fine SiC incorporation on the sinterability and mechanical properties of ZrB2″.
All the examples pinpointing the importance of thoroughly mixed small grains pertain to sintered powder compacts. Those examples: (1) illustrate the significance of using a fine-grained microstructure with well dispersed ZrB2 and SiC phases to obtain a matrix with superior mechanical properties and high oxidation resistance; (2) suggest an upper limit of about 3 μm for the size of ZrB2 grains to achieve high-performance structural materials; and (3) show that very high temperatures are required to sinter ZrB2+SiC composites (900° C. hot-pressing was used in the previous examples).
Similar results have been reported for pressureless sintering of ZrB2 using <4-wt % boron carbide (B4C) and carbon as sintering aids. W. G. Fahrenholtz et al. showed that pressureless sintering of 2-μm ZrB2 particles could be achieved at 1900° C., only if B4C and/or C were added to react with and remove the surface oxide on the particles. The oxygen-based impurities on the ZrB2 particles poisoned sintering, but the sintering temperature was reduced from 1900 to 1850° C. by reduction of the particle size to 0.5 μm and removal of oxygen content by chemical reaction; see W. G. Fahrenholtz et al., J. Am. Ceram. Soc., 91[5], 1398-1404 (2008).
A lower limit for grain size to achieve high strength has not been established. Amorphous character has been shown to be important for wear resistance of high-strength, high-fracture-toughness, and unidirectionally aligned silicon nitride (UA-SN). Studies by N. Nakamura et al. (J. Am. Ceram. Soc., 87(6), 1167-1169 (2004)) showed that ion implantation of the UA-SN surface with B+, N+, Si+ and Ti+ ions lowered the wear rate by 20 fold in the direction parallel to the grain alignment. Their transmission electron microscopy analysis of cross-sections “indicates the high wear resistance was attributed to the amouphous surface caused by the ion implantation.”*****
CFRCs obtained from infiltration of the preform with a suspension of RMC particles in a preceramic polymer typically are not high quality, because the fiber preform acts like a filter which separates the powder from the preceramic polymer. That separation leads to build up of the powder on the outside of the yarns of the preform, which in turn results in little penetration of the powder into the preform and the yarns. Regardless of the extent of penetration, the finess of the domain size is limited by the size of the particles and their agglomerates in suspension. Even when nano-sized powder is used, agglomeration of the nano particles can prevent their diffusion into the preform and the fiber threads. Only one prior effort has come close to synthesis of refractory metal-boride-SiC or refractory metal-boride+carbideSiC ceramic phases by means without addition of at least one ceramic powder. M. M. Guron et al. reported attempts to form ZrB2/ZrC/SiC and HfB2/HfC/SiC ceramic composites by pyrolysis of mixtures of Zr or Hf metal powder dispersed into blends of preceramic polymer precursors to SiC and B4C (polymethylcarbosilane and poly(norbornenyldecaborane), respectively). Although the pyrolysis of the blends of preceramic polymers without metal added gave carbon-containing SiC and B4C ceramic phases apparent by XRD of 1300 and 1600° C. chars, respectively (i.e, SiC/B4C/C composite phase after treatment at 1600° C.), pyrolysis at 1600° C. of blends containing Hf or Zr powder gave only diffraction evidence for HfB2 or ZrB2 and initial crystallization of HfC or ZrC. No SiC diffraction peaks were observed for chars derived from metals dispersed in the polymer blends. They reported that the char material was “composed of metal-boride and metal-carbide crystallites imbedded in an amorphous silicon carbon framework,” and “that ZrC and HfC crystallization was incomplete at these temperatures;” see M M. Guron, M. J. Kim, L. G. Sneddon, “A Simple Polymeric Precursor Strategy for the Synthesis of Complex Zirconium and Hafnium-Based Ultra High-temperature Silicon-Carbide Composite Ceramics,” J. Am. Ceram. Soc., 91(5), 2008, 1412-1415). In fact CFRC composites have not been formed using that precursor system, presumably due to the challenge of incorporating the Zr or Hf metal powder/preceramic polymer mixture, which is a solid, into a fiber preform.
Traditionally, ceramic bodies have been formed by sintering of powder compacts, known as the green state, which are produced by compaction of crystalline ceramic powders under high pressures. High-quality, high-performance CFRCs cannot be formed by traditional powder pressing methods and their variations, because the high pressures used damage and fracture the reinforcement fibers. The problem can be side-stepped by incorporation of very short fibers, whiskers, or chopped fibers into powder compacts. However, the short fibers do not impart the level of ultimate strength, toughness and resistance to fracturing that is desired for many commercial and industrial applications.
CFRCs comprising a nominal SiC matrix have been developed using pyrolysis of liquid or polymer based precursors for formation of the ceramic matrix within, and densification of, a fiber preform. That approach allows build-up of matrix and densification of the preform by the polymer-infiltration-pyrolysis (PIP) process without applying unwanted force on the fibers. The process relies on infiltration or flooding of the fiber preform with a solution or polymer of sufficiently low viscosity to allow penetration not only into the space between the threads of fiber but also into the microscopic spaces inside the threads of the fiber (i.e. the space between the fibrils that comprise the fiber). Polymer precursors are available for formation of SiC and SiNC matrices by the PIP process. Commercial polymers for manufacturing CFRCs using that process include Kion's Ceraset Polysilazane 20 and HTT1800 products, and Starfire's SMP 10 (allyl-hydropolycarbosilane). The CFRC composites formed by this method are often referred to as C/SiC or SiC/SiC depending on whether the reinforcement is carbon fiber or SiC fiber, respectively.
CFRC composites containing a mixed ceramic phase, such as ZrB2 intimately mixed with SiC (≦1 μm domains), are desired for applications where the composite must survive in extreme environments like hypersonic flight. That is because research has shown that the mixed ceramic phase forms a durable protective scale under those conditions, which provides more oxidation resistance than the scale of SiC alone. Those characteristics were demonstrated using monolithic pieces of the ZrB2—SiC matrix formed by hot pressing of a mixture of the powders; see M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, J. of Mat. Sci., 39, 5887-5904 (2004). Furthermore, silicon carbide composite composed mainly of SiC containing ZrB2 in an amount up to 50 volume % by sintering of powders resist penetration of, and erosion by, molten metals, such as molten steel, in a superior fashion. (J. Sugawara and S. Hosokawa, U.S. Pat. No. 4,735,923) The performance of the mixed phase ceramic is superior to that of either ceramic component alone, as SiC tends to react with molten ferrous metals and oxides thereof.